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Controlled Epitaxial Growth of Body-Centered Cubic and FaceCentered Cubic Cu on MgO for Integration on Si F. Wu* and J. Narayan Department of Materials Science & Engineering, North Carolina State University, Raleigh, North Carolina 27695, United States

ABSTRACT: The Cu/MgO interface plays a crucial role in applications. Face-centered-cubic (fcc) Cu has been reported to grow on MgO substrate (rock salt structure). However, no body-centered-cubic/tetragonal (bcc(t)) Cu has been stabilized on MgO. The special atomic structure of the bcc(t)/rock salt interface contributes to superior thermal, mechanical, and electrical properties. We report, for the first time, the epitaxial growth of bcc(t) and fcc Cu on Si(100) and Si(111) substrates using MgO(100)/TiN(100) and MgO(111)/TiN(111) buffer layers by pulsed laser deposition. We find that the deposition temperature determines the structure of Cu. At high temperature, only fcc Cu grows on both MgO/TiN(100) and MgO/ TiN(111) templates. At room temperature, an epitaxial layer of bcc(t) Cu grows pseudomorphically on a MgO(100) template up to the critical thickness, while on a MgO/TiN(111) template, the majority of Cu is fcc, and bcc(t) Cu exists occasionally in a three-dimensional island shape. The growth of these heterostructures involves epitaxy across the misfit scale by matching MgO{200} planes with bcc(t) Cu{110} planes. The integration of Cu/MgO on the technologically important Si substrate holds tremendous promise, because the novel bcc(t) Cu/MgO structure can be integrated with present-day microelectronic or nanoelectronic devices. A MgO substrate is not suitable for practical applications26 due to its high single-crystal cost, limited wafer size, and poor mechanical properties. By successfully fabricating an epitaxial MgO template on a Si substrate using TiN as the buffer layer, we can fabricate the Cu/MgO stack for integration on the technologically important silicon substrate, which is inexpensive, readily available, and widely used in present-day microelectronic devices. This holds tremendous promise because the novel bcc(t) Cu/MgO structure can be integrated with present-day microelectronic or nanoelectronic devices. All of the previously reported bcc(t) Cu were grown pseudomorphically on the substrates, either by matching the lattice19 or by matching of the sides of a unit cell of bcc(t) Cu and a primitive square mesh on {001} planes of fcc substrates.20,22 However, if bcc(t) Cu can be fabricated in a controlled way directly onto MgO(100) and MgO(111), domain matching epitaxy (DME)27−29 should be considered along with the traditional lattice matching epitaxy (LME) due to the large lattice misfit (14.2%) between these two systems.

I. INTRODUCTION Cu is the standard interconnection material in integrated circuits due to its superior electromigration resistance and lower bulk resistivity1−3 compared with Al metallization.4−6 However, due to the rapid diffusion of Cu into silicon7−9 and the consequent degradation of the semiconductor devices, a barrier layer is needed.1 MgO and TiN have been successfully utilized as good diffusion barriers for Cu for a long time.7,10,11 Furthermore, the Cu/MgO interface is of great importance because it plays a crucial role in applications such as microelectronics packaging, coating, and corrosion protection,12,13 catalysis, metal−matrix composites, recording media, etc.14 The epitaxial growth of face-centered-cubic (fcc) Cu on MgO has been reported at different temperatures.15−18 However, despite the growth of bcc Cu on substrates such as Nb,19 Ag,20 and Pd,21,22 no body-centered-cubic/tetragonal (bcc(t)) Cu has been fabricated onto MgO, which has a rock salt structure. The special atomic structure of the fcc/bcc(t) interface is believed to contribute to superior thermal, mechanical, and electrical properties.19,23−25 Consequently, the generation of bcc(t)/rock salt interfaces between bcc(t) Cu and MgO is of great scientific significance, in addition to being important for technological applications. © XXXX American Chemical Society

Received: August 6, 2013 Revised: September 13, 2013

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Figure 1. Schematic illustration of (a) 3-D atomic structure demonstrating the crystallographic relationship between bcc(t) Cu/MgO(100), (b) 2-D atomic structure of bcc(t) Cu/MgO(100), viewed along [001]bcc(t)Cu∥[001]MgO, perpendicular to the interface plane, (c) 2-D atomic structure of bcc(t) Cu/MgO(100), viewed parallel to the interface plane, (d) 3-D atomic structure demonstrating the crystallographic relationship between bcc(t) Cu/MgO(111), (e) 2-D atomic structure of bcc(t) Cu/MgO(111), viewed along [011]bcc(t)Cu∥[1̅11]MgO, perpendicular to the interface plane, and (f) 2-D atomic structure of bcc(t) Cu/MgO(111), viewed parallel to the interface plane.

For bcc(t) Cu/MgO(100), the interface planes are {100} planes for both Cu and MgO; thus the sides of a unit cell of bcc(t) Cu will match with the sides of a primitive square mesh on MgO{100}. For bcc(t) Cu/MgO(111), closed-packed atomic layers from both sides will be the interface planes for them.30 This interface orientation relationship is in accord with the reported bcc/fcc systems24,31−33 and can be explained by the Kurdjumov−Sachs34 (K−S) or Nishiyama−Wassermann (N−W) orientation relationship. In both relationships, the common interface plane is {111}fcc∥{011}bcc.35 However, for the K−S relationship, the perpendicular matching axes on the interfacial plane are [101̅]fcc∥[111̅]bcc and [1̅1̅2]fcc∥[1̅12]bcc, while those for the N−W relationship are [11̅0]fcc∥[001]bcc and [112̅]fcc∥[011̅]bcc. In this paper, we report the controlled epitaxial growth of bcc(t) and fcc Cu on both MgO(100)/TiN(100)/Si(100) and MgO(111)/TiN(111)/Si(111) by pulsed laser deposition (PLD). At high temperature, only fcc Cu grows on both MgO/TiN(100) and MgO/TiN(111) templates. At room temperature, bcc(t) Cu on MgO/TiN(100) template forms an epitaxial layer and grows pseudomorphically up to the critical thickness, corresponding to the transition where Gibbs free energy exceeds the strain free energy. However, the majority of

Cu growing on a MgO/TiN (111) template is in the stable structure (fcc), and bcc(t) Cu will grow occasionally in a threedimensional island. For both MgO/TiN(100) and MgO/ TiN(111) templates, the growth of these heterostructures involves epitaxy across the misfit scale by matching MgO{200} planes with bcc(t) Cu{110} planes. To date, this is the first report in which bcc(t) Cu is obtained on fcc-structured MgO and is the first systematic study on bcc(t) Cu on MgO templates with different orientations (MgO(100) and MgO(111)). This is also the first integration of Cu/MgO on the technologically important Si substrate. The orientation relationships between different phases were studied carefully by selected area diffraction patterns (SADP) and high-resolution transmission electron microscopy (HRTEM).

II. EXPERIMENTAL PROCEDURES Pulsed laser deposition (PLD) of each layer in the stack (Cu/MgO/ TiN/Si) was carried out in a multitarget stainless-steel chamber using a pulsed KrF excimer laser (wavelength 248 nm, pulse duration 25 ns, repetition rate 10 Hz). Details about PLD can be found in earlier publications.36,37 The deposition of TiN film was done at 625 °C in vacuum (1× 10−6 Torr). After TiN deposition, the first few monolayers (for about 500 pulses) of MgO were deposited under vacuum (1× 10−6 Torr) at 575 °C. The remaining MgO was deposited B

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Figure 2. (a) Typical TEM bright-field cross-section image showing the whole stack Cu/MgO(100)/TiN(100)/Si(100). Pt layer was deposited for protection during the FIB sample preparation process. (b) [110] zone-axis-pattern (ZAP) of Si(100)/TiN(100) interface. (C) [110] ZAP of MgO(100)/TiN(100) interface. (d) A typical HRTEM image showing the Cu/MgO(100) interface. (e) A typical HRTEM image showing the atomic structure of bct Cu at the Cu/MgO(100) interface region. at the same temperature with an oxygen pressure of 6 × 10−4 Torr. The deposition of Cu was carried out either at room temperature or at ∼300 °C. The deposition temperature determines the structure of Cu layer. The stable (fcc) structure of Cu is expected at high temperature because sufficient thermal energy is provided, while unstable bcc(t) structure of Cu is more likely to form at the Cu/MgO interface during room-temperature deposition. The microstructural characterization of the as-deposited films was performed by transmission electron microscopy (TEM) and HRTEM.36 TEM samples were prepared by the focused ion beam (FIB) technique.

that the standard bcc unit cell will be elongated along x and y axes and compressed along the z-axis. Figure 1c is the 2-D atomic-structure model viewed from MgO[110]/bcc(t) Cu[100], parallel to the interface plane. The MgO{111} planes connect with bcc(t) Cu{110} planes, and the matching plane at the interface is MgO{200}/bcc(t) Cu{110}. For comparison, the 3-D and 2-D atomic-structure models for the bcc(t) Cu/MgO(111) system are shown on the righthand side in Figure 1. Figure 1d is the 3-D atomic-structure model demonstrating the crystallographic relationship between MgO and bcc(t) Cu. The interface plane is the close-packed plane for both materials (MgO{111} and bcc(t) Cu{110}), and the close-packed direction of MgO becomes the cube edge of bcc(t) Cu. Figure 1e is the 2-D atomic-structure model viewed from ⟨011⟩ b c c ( t ) C u ∥⟨111⟩ M g O , perpendicular to the {110}bcc(t)Cu∥{111}MgO interface plane. The {110} planes of bcc(t) Cu will be elongated along ⟨011⟩ and ⟨100⟩ directions to match with ⟨112⟩MgO and ⟨110⟩MgO, respectively, corresponding to the N−W orientation relationship. Figure 1f is the 2-D atomic-structure model viewed along [001]bcc(t)Cu∥[110]MgO, parallel to the interface plane. The MgO{002} planes connect with bcc(t) Cu{110} planes, such that planar spacing matching is accomplished. The planar spacing misfit between these two materials is calculated to be (d(002)MgO − d(110)bccCu)/d(002)MgO = 3.19%, much smaller than the lattice misfit (∼31.5%) between them.

III. RESULTS AND DISCUSSION To better understand the domain matching mechanism of bcc(t) Cu on a MgO buffer layer, both 3-D and 2-D atomicstructure models (viewed either parallel or perpendicular to the interface plane) are shown in Figure 1. The phase with red balls stands for bcc(t) Cu, while the green phase stands for MgO. For the bcc(t) Cu/MgO(100) system, the sides of a unit cell of bcc(t) Cu match with the sides of a primitive square mesh on the MgO(001) plane. The 3-D atomic-structure model demonstrating this crystallographic relationship is shown in Figure 1a, in which the unit cells of bcc(t) Cu and MgO are labeled and the interface plane is bcc(t) Cu{100}/MgO{100}. Figure 1b is the 2-D atomic-structure model, and the viewing direction is [001]bcc(t)Cu∥[001]MgO, perpendicular to the interface plane. Due to the 3% misfit between abccCu and aMgO/√2, bcc(t) Cu will grow pseudomorphically on MgO so C

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Figure 3. (a) Typical TEM bright-field cross-section image showing the whole stack Cu/MgO(111)/TiN(111)/Si(111). Pt layer was deposited for protection during the FIB sample preparation process. (b) [110] zone-axis-pattern (ZAP) of Si(111)/TiN(111) interface. (c) [110] ZAP of MgO(111)/TiN(111) interface. (d−f) HRTEM images of fcc Cu/MgO (111) interface at different magnifications, viewed along [110] direction. (g) Schematic illustration of the hexagonal dislocation network on (111) planes between two close-packed lattices (fcc Cu/MgO) with a cube-on-cube relationship.

Figure 2a is a typical TEM bright-field cross-section image showing the whole stack Cu/MgO(100)/TiN(100)/Si(100). The Pt layer was deposited onto this stack for protection during the FIB sample preparation process. Individual Cu islands grow on the MgO(100) template, indicating a 3-D growth mode of Cu. The thicknesses of TiN, MgO, and Cu are about 136, 150, and 40 nm, respectively. To study the microstructures of the buffer layers and the orientation relationships between them, selected area diffraction patterns (SADP) were obtained at each interface, as shown in Figure 2b,c. Figure 2b shows the [110]

zone-axis pattern of Si/TiN interface, clearly demonstrating a cube-on-cube relationship between the TiN and Si substrate. Both {002} and {111} planes of TiN and Si align parallel to each other. The diffraction points corresponding to the same planes of TiN and Si split but appear along the same directions due to the lattice misfit of 24.6% (aSi = 5.43 Å; aTiN= 4.24 Å38). The interface is TiN{001}/ Si{001}. The domain matching epitaxy (DME)27−29 between single-crystal TiN and Si substrate has been illustrated elsewhere.28 Figure 2c shows the [110] zone-axis-pattern of MgO/TiN interface. The D

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Figure 4. (a) A typical HRTEM image in which island-shaped bcc(t) Cu and epitaxial fcc Cu coexist on MgO(111). (b) Enlarged HRTEM image showing the island region where bcc Cu lattice appears. (c) [110] ZAP of fcc Cu. (d) [110] ZAP of MgO. (e) [111] ZAP of bcc Cu. (f) Simulated [110] ZAP of fcc Cu and [111] ZAP of bcc Cu.

template. From Figure 3e,f, it can be seen that the fcc Cu/ MgO(111) interface is atomically flat and sharp, without any secondary phase over a large area. Periodic contrast is visible along the interface due to the existence of misfit dislocations with a certain separation, indicating that the fcc Cu/MgO(111) interface is semicoherent. The fcc Cu{002} planes connect with MgO{002} planes across the interface between the interfacial dislocations and bend only within the localized width of the misfit dislocations. Since Cu has a large lattice misfit of 14.2% with respect to MgO, their matching should be explained by the DME27−29 paradigm instead of the traditional LME, in which the misfit between the matching materials is typically less than 7−8%.29 Two domains, in which seven planes of fcc Cu{002} match with six planes of MgO{002} and six planes of fcc Cu{002} match with five planes of MgO{002}, will alternate with a certain frequency. The domain matching mechanisms are illustrated in Figure 3e,f. The highly magnified HRTEM image shown in Figure 3f demonstrates a periodic contrast variation along the [1̅12] direction. The misfit dislocations lying on the interfacial plane are clearly seen with an average separation of 1.53 ± 0.10 nm. The Burgers vector of the interfacial dislocations can be 1/ 2⟨101̅⟩, which is a perfect lattice vector of both MgO and Cu. According to Bollmann’s O-lattice theory,40 a hexagonal dislocation network, which is illustrated in Figure 3g, will exist on (111) planes between two close-packed lattices with a cube-on-cube relationship. The misfit dislocations are lying along ⟨112⟩ directions, with Burgers vectors of 1/2⟨110⟩, such that all of them are edge dislocations, which are the most effective dislocations accommodating the lattice mismatch.40,41

overlapping of diffraction points along all three directions (due to their close lattice constants (aMgO = 4.216 Å)) indicates clearly a cube-on-cube relationship between MgO and TiN. The interface is TiN{001}/MgO{001}. These diffraction patterns demonstrate the epitaxial growth of MgO(100)∥TiN(100)∥Si(100). To investigate the atomic structure between Cu and MgO(100) interface, we collected HRTEM images, as shown in Figure 2d. First, an epitaxial layer of bcc(t) Cu grows pseudomorphically on MgO(100), then after it reaches the critical thickness, the Gibbs free energy exceeds the strain free energy and the unstable bcc(t) Cu phase transforms to the stable fcc structure. The atomic structure at the interface region is enlarged in Figure 2e. The viewing direction is MgO[110]/ bcc(t) Cu[100]. The bcc(t) Cu is 7 ± 1 monolayers thick, and the bcc(t) Cu{011} planes connect with MgO{111} planes across the interface. The matching planes are MgO{200}/ bcc(t) Cu{110}, in accord with the matching relationship between bcc Ni and MgO.39 For comparison, the Cu/MgO(111)/TiN(111)/Si(111) heterostructure was also investigated with TEM. Figure 3a shows the whole stack, in which MgO and TiN layers are measured to be ∼150 nm thick, and the Cu layer is thinner than 50 nm. SADP of Si/TiN and MgO/TiN interfaces shown in Figure 3b,c demonstrate the epitaxial growth of MgO(111)∥TiN(111)∥Si(111). Figure 3d,e,f are HRTEM images showing the fcc Cu/MgO(111) interface at different magnifications, and all of them are viewed along the [110] direction. Unlike the situation for the Cu/MgO(100) interface, the majority of the Cu phase remains fcc-structured on the MgO/TiN(111) E

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applicable to wavy (curved) interfaces as well. Wei et al.46 reported that the wavy interface gave rise to a variety of orientation relationships and the interface structure may be more complex and involve disconnections. At the top/bottom layers, K−S34 and N−W orientations (common interface plane, {111}fcc∥{011}bcc35) dominated between two adjacent (V/Ag) layers, while on the slopes of the wavy interface, the Bain orientation relationship (common interface plane {100}fcc∥{100}bcc47) was preferred. However, although the orientation relationships between two adjacent layers may change from one position to another for wavy interfaces, the matching planes for bcc(t) Cu, fcc Cu, and MgO phases will not change, which are bcc(t) Cu{110}∥MgO{200}∥fcc Cu{200} planes. The DME mechanism determines that eq 2 holds true for wavy interfaces as well. This is important because curved surfaces play a critical role in dislocation nucleation as a result of interfacial steps and strain relaxation.48,49 The simulated diffraction patterns for fcc Cu and bcc(t) Cu are shown together in Figure 4f, which corresponds well to the zone axis patterns we obtained for the two phases.

The spacing between the dislocation lines can be calculated according to d = aMgOaCu/[1.414(aMgO − aCu)].41 Taking the projection effect into consideration, the dislocation spacing viewed along [110] direction should be 1.56 nm for the Cu/ MgO(111) interface, which matches with our observations. Although the majority of the Cu phase growing on the MgO(111) template remains fcc-structured, 3-D island shaped bcc Cu grows on MgO(111) occasionally as well. Figure 4a is a typical HRTEM image in which bcc(t) and fcc Cu coexist on MgO(111). The Cu/MgO(111) interface is drawn as a black line as a guide to the eyes. As easily seen from Figure 4a, the majority of the Cu phase remains fcc-structured and the MgO{002} planes match with Cu{002} planes as illustrated previously. However, within the red-circle region, bcc(t) Cu forms directly on MgO(111). After several monolayers of epitaxial growth, the unstable bcc-structured Cu turns back to its equilibrium structure again. The bcc(t) to fcc phase transformation of Cu can be understood as the competition between the strain energy and the Gibbs free energy associated with this phase transformation. When the bcc(t) phase grows beyond the critical thickness, the strained layer will relax to its bulk structure by introduction of misfit dislocations. The region in which bcc(t) Cu lattice appears is enlarged as shown in Figure 4b. The critical thickness can be determined from eq 1:27,42,43 2μ(1 + υ)εr 2(tc) ≤ ΔGv (1 − υ)

IV. CONCLUSIONS Bcc(t) and fcc Cu were fabricated onto MgO(100)/TiN(100)/ Si(100) and MgO(111)/TiN(111)/Si(111) for integration on silicon substrates. At high temperature, only fcc Cu grows on both MgO(100) and MgO(111) templates. At room temperature, an epitaxial layer of bcc(t) Cu will first grow on MgO(100) pseudomorphically and then transform back to fcc Cu after exceeding the critical thickness. In contrast, the majority of Cu on MgO(111) would remain fcc-structured, and island-shaped bcc(t) Cu will occasionally grow on MgO(111). For both MgO/TiN(100) and MgO/TiN(111) templates, the growth of these heterostructures involves epitaxy across the misfit scale by matching MgO{200} with bcc(t) Cu{110} planes. The morphology of the Cu layer can be controlled by the orientation of MgO templates and the deposition temperatures.

(1)

where εr is the residual strain, which has a thickness dependence, μ is the shear modulus, υ is the Poisson’s ratio, and ΔGv is the Gibbs free energy difference between fcc and bcc(t) Cu. This model takes into account the strain distribution in the new phase, which is not accounted by the earlier thermodynamic model.44 The neglect of the effect of coherency strains will result in a large error, due to the large strain within the bcc(t) Cu/MgO heterostructure. To better understand the orientation relationship between each phase, the selected area diffraction patterns of fcc Cu, MgO, and bcc(t) Cu are demonstrated in Figure 4c,d,e, respectively. Figure 4c,d shows [110] zone axis patterns for fcc Cu and MgO, respectively, from which it is clearly seen that the interface plane is fcc Cu{1̅11}/MgO{11̅ 1}. The {002} lattice planes of Cu match with those of MgO across the {1̅11} interface plane. Figure 4e is the [111] zone axis pattern for bcc(t) Cu, which indicates that the interface plane is actually the {10̅ 1} lattice plane for bcc(t) Cu. The {002} lattice planes of MgO now connect {011̅} lattice planes of bcc(t) Cu across the interface plane. Using MgO (a = 4.216 Å) as the internal calibration standard, the lattice parameter of bcc(t) Cu is calculated to be 2.886 ± 0.010 Å, which fits well to the reported range of the lattice constant of bulk bcc(t) Cu (from 2.87 to 2.96 Å45). Compared with small lattice misfit systems (∼0.3%) in which bcc(t) Cu grows pseudomorphically on substrates such as Ag,20 planar spacing matching instead of lattice matching holds true for bcc(t) Cu/MgO. The following equation satisfies the domain matching among these three phases: aMgO a a = fccCu = bccCu (2) 2 3 2

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AUTHOR INFORMATION

Corresponding Author

*Mailing address: 708f Chappell Drive, Raleigh, NC, USA. Phone 919-760-0269. E-mail: [email protected]. Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS F.W. and J.N. acknowledge the support by the National Science Foundation of the United States (Grant No. DMR41104667). The authors thank Dr. Srinivasa Rao Singamaneni for proofreading this manuscript and providing technical support.

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REFERENCES

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It should be pointed out that although this relationship is derived from flat interfaces presented in this paper, eq 2 is F

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dx.doi.org/10.1021/cg4011983 | Cryst. Growth Des. XXXX, XXX, XXX−XXX